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Abstract

InAs/GaAs(001) quantum dots grown by droplet epitaxy were investigated using electron
microscopy. Misfit dislocations in relaxed InAs/GaAs(001) islands were found to be
located approximately 2 nm above the crystalline sample surface, which provides an
impression that the misfit dislocations did not form at the island/substrate interface.
However, detailed microscopy data analysis indicates that the observation is in fact
an artefact caused by the surface oxidation of the material that resulted in substrate
surface moving down about 2 nm. As such, caution is needed in explaining the observed
interfacial structure.

Keywords:

Background

Semiconductor quantum dots (QDs) are attracting increasing research interest because
of their important potential applications in electronic and optoelectronic devices
[1-4]. Among the many techniques to produce QDs, coherent island formation through heteroepitaxial
growth of semiconductor materials has been the most important technique because of
the possible combination of the QD growth and semiconductor integration techniques.
Two different methods have been used to grow epitaxial QDs. The first one is the classical
Stranski-Krastanow (S-K) growth, [5,6] in which a semiconductor material with larger lattice parameters is first deposited
on the substrate with smaller lattice parameters layer-by-layer, forming a wetting
layer, followed by island formation to partially release the strain energy caused
by the lattice mismatch between the epilayer and the substrate. The other method is
droplet epitaxy [7-9] that has been used for the growth of III-V semiconductor QDs [10,11] by firstly introducing liquid nanodroplets of the group III element on the substrate
and then exposing the droplets to a gas-phase flow of the group V element. Different
from the S-K growth mode, droplet epitaxy does not depend on lattice mismatch and
therefore can be applied to more materials systems.

For both the S-K growth and the droplet epitaxy, misfit dislocations will be introduced
to further release the strain caused by the lattice mismatch when the size of a QD
reaches a critical value [12,13]. Because misfit dislocations produce deleterious effects on QD properties, the mechanisms
on how misfit dislocations are generated have been extensively investigated both experimentally
and theoretically [11-20]. Based on the fact that the highest stress in a QD occurs at the corner where the
island meets the substrate, it has been suggested that strain-relieving perfect misfit
dislocations are generated at the island edge, when the island reaches its critical
size [14-16]. The generation and morphology of perfect misfit dislocations have been considered
theoretically [17-19]. Partial misfit dislocations have also been identified in relaxed islands [9,20] and are believed to be energetically more favourable than perfect misfit dislocations
in some regions of QDs [21]. In some situations, e.g. Ge/Si(001) QDs grown at high temperature, partial misfit
dislocations can originate from the surface of QD islands and then glide to the island/substrate
interface [22]. While most misfit dislocations are located at the island/substrate interface, it
is interesting to note an exception in a recent report that presented evidence of
misfit dislocations located above the island/substrate interface in an S-K-grown GaSb/GaAs(001)
system [23]. The phenomenon was explained based on the compressive stress induced on the GaSb
islands by the GaAs substrate due to lattice mismatch between the epilayer and the
substrate [23]. In this letter, we conducted electron microscopy characterization of InAs/GaAs(001)
QDs grown by droplet epitaxy. Misfit dislocations in relaxed InAs/GaAs(001) islands
were found to be located about 2 nm above the island/substrate interface. However,
detailed analysis of the transmission electron microscopy (TEM) data indicates that
the observation is in fact an artefact caused by surface oxidation of the material
that resulted in substrate surface moving down about 2 nm. We therefore conclude that
caution is needed when explaining the interfacial structure of the QDs.

Methods

A double-layer InAs/GaAs(001) QD sample was grown by droplet epitaxy using metal organic
vapour phase epitaxy. The QD sample was grown on a semi-insulating GaAs(001) substrate
in a horizontal flow reactor (AIX200/4, AIXTRON SE, Herzogenrath, Germany) at a pressure
of 100 mbar. Trimethylindium, trimethylgallium and AsH3 were used as the precursors and ultrahigh-purity H2 as the carrier gas. Firstly, a 200-nm GaAs buffer layer was deposited at 650°C, then
the temperature was reduced to 500°C and the growth was interrupted for 10 s with
all sources removed from the reactor to eliminate the influence of AsH3 source on the later deposition of indium droplets. After that, two monolayers of
indium droplets were deposited, which were then immediately exposed to the AsH3 flow (3.0 × 10−4 mol/min) for 15 s. The InAs QDs were capped immediately by a 10-nm GaAs capping layer
at 500°C. Then, the growth temperature was ramped up to 650°C and a 100-nm GaAs layer
was deposited. The InAs QDs were capped immediately by a 100-nm GaAs capping layer
whilst the temperature was ramped up to 650°C. Finally, a surface layer of InAs QDs
was grown in the same fashion as the buried InAs QDs.

Only the structures of QDs on the top layer were investigated. Cross-sectional TEM
specimens were prepared using a Gatan precision ion polishing system (Gatan, Inc.,
Pleasanton, CA, USA) with Ar+ energy of 2.5 keV. Structural characterization of the QDs was conducted using scanning
electron microscopy (SEM; Zeiss Ultra+, Carl Zeiss, Inc., Oberkochen, Germany) operated
at 2 kV and TEM (JEM-3000 F, JEOL Ltd., Akishima-shi, Japan) operated at 300 kV. Quantitative
compositional analysis was conducted using X-ray energy-dispersive spectroscopy (XEDS)
in JEM-2200 TEM (JEOL Ltd.) operated at 200 kV and the ESPRIT software. The electron
probe size used for the XEDS was 1 nm. XEDS data collection time was controlled to
make sure that high counts (larger than 10 thousands) were obtained for datum points
so that the statistical errors were less than 1%.

Results and discussion

Figure 1a shows a typical plan-view SEM image of the QDs on the surface. Large islands with
widths (measured along the [110] direction) larger than approximately 30 nm are usually
elongated and faceted, showing typical strain-relaxed morphologies [24]. Most of small islands with widths smaller than approximately 15 nm have an elliptical
shape, although a small number of them are circular. The large islands are all relaxed
through the formation of misfit dislocations. Figure 1b shows a typical cross-sectional high-resolution TEM image of a relaxed large island on the sample
surface. Three misfit dislocations, which are indicated with three white arrows, are
seen at the same atomic layer in the island, which is approximately 2 nm above the
crystalline surface of the substrate. A white line marked as ‘1’ indicates the crystalline
surface of the substrate, and the other white line marked with ‘2’ specifies the position
of the atomic layer on which the misfit dislocations lie.

Figure 1.Morphology of the QDs. (a) An SEM image of the top surface of the InAs/GaAs(001) QDs; (b) a [] cross-sectional TEM image showing a relaxed InAs island on the GaAs(001) substrate.
Line 1 and line 2 indicate the positions of the crystalline surface of the substrate
and the surface of the dark amorphous layer, respectively. Three misfit dislocations
in the relaxed island are marked with three white arrows.

The phenomenon of misfit dislocations positioned above the crystalline substrate surface
seen in Figure 1 is very similar to the phenomenon reported in [23]. However, another interesting phenomenon is also seen in Figure 1. There are two amorphous structures with distinctly different contrasts above the
crystalline structure: one is a thin amorphous layer with a thickness of about 2 nm
immediately above the crystalline substrate and island surface and with contrast clearly
darker than the other part of the amorphous area above the sample. The top surface
of the 2-nm thin amorphous layer is located at exactly the same level as the atomic
plane where misfit dislocations lie, i.e. the layer indicated by line 2. While the
amorphous structure with bright contrast is the epoxy used for cross-sectional sample
preparation, it is not clear what the dark amorphous structure is.

XEDS analysis was conducted to explore the composition of the dark-contrast amorphous
layer and surrounding areas. Figure 2 shows the results obtained from an island and its nearby area. Line scans were conducted
along lines ABCD and EFG shown in Figure 2a with the interval of the datum points at approximately 2 nm. O, Ga, As and In were
detected in the material. The quantitative concentrations of these elements as a function
of positions along lines ABCD and EFG are shown in Figure 2b,c, respectively. From Figure 2b, a small amount of In content is detected in the substrate; this is caused by the
combined effect of the contamination induced during the sample preparation process
and the electron beam spreading to the island through the scatter of the electron
beam in the TEM sample. The latter effect can be removed when XEDS was conducted in
the substrate far away from the island, as shown in Figure 2c. Another possible reason that In is detected in the substrate is the mutual diffusion
of In into the substrate and Ga into islands during the high-temperature sample deposition
process. This explains the phenomenon that over 10% of Ga is detected in the island.
The inter-diffusion process of In and Ga is in fact an effective way to reduce the
strain energy of the system [16,25,26].

Figure 2.XEDS analysis. (a) A [] cross-sectional TEM image of a relaxed InAs island. Straight lines ABCD and EFG
indicate the positions from which XEDS data were obtained; (b) XEDS data detected along the line ABCD; (c) XEDS data detected along the line EFG.

The most interesting phenomena revealed by the XEDS data in Figure 2 are (1) although O is detected everywhere in the sample, which is caused by surface
oxidation after TEM sample preparation, the O content increases significantly at the
dark amorphous layer; (2) the contents of Ga at the position immediately below misfit
dislocations and at the dark-contrast amorphous layer immediately above the crystalline
substrate are very high; and (3) the In content at the dark-contrast amorphous layer
covering the island is very high. All these evidence point to the conclusion that
the dark-contrast amorphous layer is actually an oxidised layer that was originally
the surface of the crystalline GaAs substrate and the surface of InAs islands. The
oxidised amorphous layer formed after the epitaxial sample surface was exposed to
the air and therefore was only seen on the sample surface, not in the buried QD layer.
Therefore, misfit dislocations in large relaxed QD islands formed exactly at the interface
of the epilayer and the substrate, not at a level approximately 2 nm above the interface
as it looks in Figure 1. Because oxidation of the surface of semiconductor nanostructures has been widely
reported in literatures [27-29], not just in the InAs/GaAs QD structure reported here, caution is needed when investigating
the interfacial structures of epitaxial materials.

Conclusions

In summary, misfit dislocations in InAs/GaAs(001) QDs grown by droplet epitaxy are
observed to be located approximately 2 nm above the crystalline substrate surface.
However, detailed compositional analysis suggests that this is an artefact caused
by surface oxidation. The oxidised surface is of an amorphous structure with a thickness
of approximately 2 nm.

Competing interests

The authors declare that they have no competing interests.

Authors’ contributions

ZBC, WL and XZL designed the study. ZBC conducted the microscopy experiments. WL,
HHT and CJ grew the QD specimens. ZBC and XZL wrote the paper. All authors discussed
the results and contributed to the paper. All authors read and approved the final
manuscript.

Authors’ information

ZBC is a postgraduate student. Dr WL is a research associate professor. Dr BC is a
postdoctoral research associate. Dr YBW is an Australian Research Council (ARC) Australian
Postdoctoral Fellow. Dr XZL is an associate professor and ARC Future Fellow. Dr HHT
is a senior fellow and ARC Future Fellow. Dr JZ is a professor and ARC Future Fellow.
SPR is the Director of the Australian Centre for Microscopy and Microanalysis, the
Executive Director and CEO of the Australian Microscopy and Microanalysis Research
Facility, and the Director of the Bandwidth Foundry International Pty Ltd. Dr CJ is
a distinguished professor and ARC Laureate Fellow.

Acknowledgements

The authors are grateful for the scientific and technical support from the Australian
Microscopy & Microanalysis Research Facility node at the University of Sydney. We
thank the Australian National Fabrication Facility for providing access to growth
facilities used in this work. This research was financially supported by the Australian
Research Council.